Improving the Mechanical Response of Al-Mg-Si 6082 Structural Alloys during High-Temperature Exposure through Dispersoid Strengthening

Abstract

The feasibility and efficacy of improving the mechanical response of Al-Mg-Si 6082 structural alloys during high temperature exposure through the incorporation of a high number of α-dispersoids in the aluminum matrix were investigated. The mechanical response of the alloys was characterized based on the instantaneous high-temperature and residual room-temperature strengths during and after isothermal exposure at various temperatures and durations. When exposed to 200 °C, the yield strength (YS) of the alloys was largely governed by β" precipitates. At 300 °C, β" transformed into coarse β', thereby leading to the degradation of the instantaneous and residual YSs of the alloys. The strength improvement by the fine and dense dispersoids became evident owing to their complementary strengthening effect. At higher exposure temperatures (350-450 °C), the further improvement of the mechanical response became much more pronounced for the alloy containing fine and dense dispersoids. Its instantaneous YS was improved by 150-180% relative to the base alloy free of dispersoids, and the residual YS was raised by 140% after being exposed to 400-450 °C for 2 h. The results demonstrate that introducing thermally stable dispersoids is a cost-effective and promising approach for improving the mechanical response of aluminum structures during high temperature exposure.

Full text

materials
Article
Improving the Mechanical Response of Al–Mg–Si
6082 Structural Alloys during High-Temperature
Exposure through Dispersoid Strengthening
Jovid Rakhmonov 1, * , Kun Liu 1, Paul Rometsch 2, Nick Parson 2and X.-Grant Chen 1, *
1Department of Applied Science, University of Quebec at Chicoutimi, Saguenay, QC G7H 2B1, Canada;
kun.liu@uqac.ca
2Arvida Research and Development Center, Rio Tinto Aluminum, Saguenay, QC G7S 4K8, Canada;
Paul.Rometsch@riotinto.com (P.R.); nick.parson@riotinto.com (N.P.)
*Correspondence: jovid.rakhmonov1@uqac.ca (J.R.); xgrant_chen@uqac.ca (X.-G.C.)
Received: 20 October 2020; Accepted: 18 November 2020; Published: 23 November 2020


Abstract:
The feasibility and efficacy of improving the mechanical response of Al–Mg–Si 6082
structural alloys during high temperature exposure through the incorporation of a high number of
α
-dispersoids in the aluminum matrix were investigated. The mechanical response of the alloys was
characterized based on the instantaneous high-temperature and residual room-temperature strengths
during and after isothermal exposure at various temperatures and durations. When exposed to
200
◦
C, the yield strength (YS) of the alloys was largely governed by
β
” precipitates. At 300
◦
C,
β
”
transformed into coarse
β
’, thereby leading to the degradation of the instantaneous and residual YSs
of the alloys. The strength improvement by the fine and dense dispersoids became evident owing to
their complementary strengthening effect. At higher exposure temperatures (350–450
◦
C), the further
improvement of the mechanical response became much more pronounced for the alloy containing
fine and dense dispersoids. Its instantaneous YS was improved by 150–180% relative to the base alloy
free of dispersoids, and the residual YS was raised by 140% after being exposed to 400–450
◦
C for
2 h. The results demonstrate that introducing thermally stable dispersoids is a cost-effective and
promising approach for improving the mechanical response of aluminum structures during high
temperature exposure.
Keywords:
Al–Mg–Si 6082 alloys; microstructure; high-temperature mechanical properties; residual
mechanical behavior; α-Al(MnFe)Si dispersoids
1. Introduction
Owing to their preferable strength-to-weight ratio, good corrosion resistance, and weldability,
Al–Mg–Si 6xxx alloys (typically 6061 and 6082 alloys) are increasingly used in load-bearing structural
applications, such as land-based vehicles, marine crafts, light rails, bridge decks, off-shore platforms,
and building structures [
1
–
6
]. Such aluminum structures can be subjected to unintentional fire
exposures; therefore, fire safety is a major concern in their design and applications [3,4,7].
To evaluate the mechanical behavior of aluminum alloys under fire conditions, different test
methods can be adopted [
3
], such as transient and steady-state mechanical tests, as well as the creep test.
Based on data from steady-state tests, codes and regulations describing the design requirements for
fire-prone aluminum structures have been established [
8
]. Owing to the lower melting temperature of
aluminum alloys compared to those of common structural metals (e.g., iron and steel), the temperature
range representative of realistic fire events and relevant for the mechanical properties is 150–450
◦
C [
4
].
In addition to the temperature, the exposure period at a given temperature is a key parameter that
Materials 2020,13, 5295; doi:10.3390/ma13225295 www.mdpi.com/journal/materials

Materials 2020,13, 5295 2 of 13
determines the resistance of aluminum structures to failure during and after fire. Most studies [
7
] and
norms [
8
] indicate that the significant exposure time is between 0.5 and 2 h at a given temperature;
0.5 h refers to the critical time for safe evacuations and 2 h to the time the structure needs to maintain
the adequate strength before the fire is extinguished [
3
,
8
]. However, in most studies, the constitutive
behavior of aluminum structural materials at various temperatures and exposure periods have
been assessed without considering the microstructural changes during the fire. Understanding the
evolution of the microstructure and mechanical properties under fire conditions is essential for tailoring
microstructures such that the aluminum structures perform better in fire events. Moreover, the number
of studies of the high-temperature mechanical behavior of Al–Mg–Si alloys under various fire exposure
severities is limited.
Al–Mg–Si 6xxx alloys are heat-treatable alloys, and their primary strengthening mechanism is the
precipitation of nanoscale
β
”/
β
’-MgSi precipitates, which significantly increase the room temperature
strength. However, these alloys experience a significant strength drop when exposed to temperatures
higher than 200
◦
C, since the
β
”/
β
’ precipitates are highly sensitive to elevated temperature [
9
–
11
].
The damage extent of the properties of aluminum alloys depends on the fire exposure conditions
(temperature and time). The coarsening and subsequent phase transformation of strengthening
β
”/
β
’
precipitates in 6xxx alloys are the main factors contributing to the degradation of the mechanical
properties at elevated temperatures [1,3,12].
Recent studies have revealed that the mechanical properties (strength and creep resistance) of
aluminum alloys at elevated temperatures can be substantially improved by introducing a high number
of thermally stable
α
-Al(MnFe)Si dispersoids (referred as
α
-dispersoids hereafter) [
13
–
17
]. The most
significant advantage of the dispersoid-strengthening over the traditional precipitation-strengthening
in 6xxx alloys is their excellent thermal stability at elevated temperatures. The
α
-dispersoids can be
formed in an
α
-Al matrix during the homogenization of Mn-containing 6082 alloys [
18
,
19
]. However,
these dispersoids are coarse (diameter ranging from 100–200 nm) and exhibit low density, which leads to
an insufficient strengthening effect [
20
]. In general, the efficiency of the dispersoid-strengthening effect
depends highly on their size and density. The optimization of homogenization conditions to induce
the precipitation of fine and dense dispersoids will be a new avenue to improve the high-temperature
mechanical properties, which increase the time to failure/collapse of the aluminum 6082 structures,
and hence improve the safety in case of a fire.
Furthermore, the post-fire (residual) mechanical behavior of aluminum structures is a key
parameter that determines the stability and integrity of the entire structure [
21
]. After the fire,
the aluminum structures (members) might be replaced, repaired, or directly reused, depending on
the severity of the fire damage [
2
,
21
]. The residual mechanical behavior of aluminum alloys after
a fire has been investigated in previous studies [
2
,
21
]. The residual yield strength (YS) of 6082-T6
alloy exposed to 450
◦
C and then cooled to room temperature has been reported to be 65 MPa, which
was approximately 20% of the original YS before the high-temperature exposure [21]. The possibility
of retaining fineness and high density of
α
-dispersoids in the microstructure of 6082 alloys during
thermomechanical processes might also provide the aluminum structures with higher residual strength
and better fire resistance, in light of the highly diminishing strengthening effect of
β
”/
β
’ precipitates
upon thermal exposure.
The aims of this study were (1) to evaluate the microstructural evolution and associated mechanical
properties of 6082 structural alloys exposed to different temperatures and times (simulated fire exposure),
and (2) to assess the feasibility and efficacy of enhancing the mechanical response of 6082 structural
alloys by the incorporation of fine and dense
α
-dispersoids, in addition to the
β
”-MgSi precipitates,
into their microstructure.
2. Materials and Methods
Two 6082 type alloys (without and with Mn) were prepared and cast in direct chill (DC)-cast billets
with diameters of 101 mm. The chemical compositions of the alloys analyzed by optical emission

Materials 2020,13, 5295 3 of 13
spectroscopy are provided in Table 1. To promote the precipitation of a large amount of dispersoids [
11
],
the DC-cast billets of the alloys were heat-treated at a relatively low temperature (400
◦
C) for 5 h before
the extrusion process [
11
]; the 0.72Mn alloy treated under this condition is referred to as “0.72Mn(L)”.
In addition, some DC-cast billets of the 0.72Mn alloy were heat-treated at 550
◦
C for 5 h, which are
hereafter referred to as “0.72Mn(H)”. This typical industrial high temperature homogenization results
in few and coarse dispersoids in the aluminum matrix [
18
,
19
]. Before the extrusion, all billets were
inductively heated to 500
◦
C and then quickly transferred to the extrusion press. The extrusion was
conducted at a ram speed of 10 mm/s to produce rods with diameters of 17.8 mm. Subsequently,
the rods exiting the die traveled through a bath with agitating water, which served as a water quenching
treatment. The rods were then subjected to an artificial aging at 180 ◦C for 5 h (T5 treatment).
Table 1. Chemical compositions of experimental alloys (wt.%).
Alloys Mg Si Mn Fe Ti Al
0Mn (base) 0.83 1.01 0 0.22 0.018 Bal.
0.72Mn 0.84 1.02 0.72 0.23 0.016 Bal.
To determine the mechanical behavior of the experimental alloys, compressive YS tests were
conducted with a Gleeble 3800 thermomechanical simulator (Dynamic Systems Inc., Austin, TX, USA)
unit at room and elevated temperatures. The mechanical properties of the three alloy variants to
the thermal exposure were characterized by steady-state mechanical testing [
3
]. Different exposure
temperatures (200, 300, 350, 400, and 450
◦
C) and typical two exposure times (0.5 h and 2 h) were
chosen in the experiment. T5-treated rods from the three alloy variants were heated to the target
exposure temperatures with a heating rate of 0.25
◦
C/s and held either for 0.5 h or 2 h, followed by
water quenching. Cylindrical specimens of 15 mm height and 10 mm diameter were machined from the
thermally exposed extrusion rods. During the Gleeble compression tests, the specimens were heated to
the same target temperature at a rate of 2
◦
C/s, and held for 180 s to ensure a uniform temperature
distribution. Subsequently, the specimens were compressed at a strain rate of 10
−3
s
−1
to reach a total
strain of 0.2. Test reproducibility was ensured by testing at least three samples per condition.
An optical microscope (OM, Nikon Instruments Inc., Tokyo, Japan), a scanning electron microscope
(SEM, JEOL, Tokyo, Japan) equipped with electron backscatter diffraction (EBSD), and a transmission
electron microscope (TEM, JEOL JEM-2100, Tokyo, Japan) were used to characterize the microstructure
of the alloys. The grain structures after the extrusion and after thermal exposure were investigated by
EBSD (samples were sectioned in the extrusion direction). The precipitates and dispersoids in the alloys
under T5 conditions and after thermal exposure were characterized with the TEM, and TEM thin foils
were obtained by electro-chemical polishing. The characteristics of the
β
” precipitates (e.g., average
volume (
Vp
), number density (
Nv
), and volume fraction (
fp
)) were quantified with the method in [
22
].
Moreover, the dispersoids were quantitatively characterized based on their average diameter (
D
),
number density (
Nv
), and volume fraction (
fd
), which were determined according to the methodology
provided in [23].
3. Results
3.1. Microstructure and Mechanical Properties under T5 Conditions
Figure 1shows the typical TEM images of the experimental alloys under T5 conditions. The results
reveal uniformly distributed, fine, needle-shaped precipitates (approximately 4 nm diameter and
30 nm length) in the aluminum matrix of all three alloy variants. According to their morphology and
size, the precipitates were identified as
β
” phase [
24
]. The quantitative results of the
β
” precipitates
are listed in Table 2. Compared to the 0Mn base alloy, the Mn-containing alloys were characterized by
higher number densities and volume fractions of the β” phase.

Materials 2020,13, 5295 4 of 13
Materials 2020, 13, x FOR PEER REVIEW 4 of 13
because the aging treatment at 180 °C had no effect on the coarsening of the dispersoids. The
dispersoids in the matrix were identified as α-Al(FeMn)Si [10,11], and their quantitative results are
presented in Table 2. The average equivalent diameter (
) and number density () of the dispersoids
in the 0.72Mn(L) alloy were measured to be 40 nm and 430 µm
−3
, respectively
(versus 
 of 146 nm
and  of 11 µm
−3
in 0.72Mn(H) alloy), while the volume fractions of the dispersoids (ʄ) in both
alloys were similar (approximately 0.9%).
Figure 1. Bright-field TEM images of precipitates and dispersoids in aluminum matrices of T5-treated
(a) 0Mn, (b) 0.72Mn(H), and (c) 0.72Mn(L) alloys.
Figure 2. TEM images showing α-Al(FeMn)Si dispersoids embedded in aluminum matrices of (a)
0.72Mn(H) and (b) 0.72Mn(L) alloys. Images were taken near [001]
Al
zone axes.
Table 2. Quantitative TEM results of β” precipitates and α-Al(FeMn)Si dispersoids in three alloys.
Alloy β” phase
α
-Al(FeMn)Si Dispersoids
, nm
3
, nm
−3
,% 
 , µm
−3
,%
0Mn 410 1.26 × 10
–5
0.52 - - -
0.72Mn(H) 459 1.40 × 10
–5
0.62 146 11 0.91
0.72Mn(L) 468 1.33 × 10
–5
0.64 40 430 0.88
Based on the EBSD results of the as-extruded materials, it could be observed that the 0Mn alloy
experienced full recrystallization during extrusion because the grains were coarse, equiaxed, and free
of the substructure (Figure 3a) [25]. A <001> recrystallization texture highly dominated in this alloy
(Figure 3a). The average equivalent diameter of the grains in 0Mn alloy were measured as ~220 µm.
Figure 1.
Bright-field TEM images of precipitates and dispersoids in aluminum matrices of T5-treated
(a) 0Mn, (b) 0.72Mn(H), and (c) 0.72Mn(L) alloys.
Table 2. Quantitative TEM results of β” precipitates and α-Al(FeMn)Si dispersoids in three alloys.
Alloy
β” Phase α-Al(FeMn)Si Dispersoids
Vp, nm3Nv, nm−3fp, % ¯
DNv,µm−3fd,%
0Mn 410 1.26 ×10−50.52 - - -
0.72Mn(H) 459 1.40 ×10−50.62 146 11 0.91
0.72Mn(L) 468 1.33 ×10−50.64 40 430 0.88
A key difference between the TEM microstructures of the various alloy conditions was the presence
of a number of dispersoids in addition to the
β
” precipitates in the 0.72Mn(H) (Figure 1b) and 0.72Mn(L)
(Figure 1c) alloys. The dispersoids in the 0.72Mn(L) alloy were much finer and denser than those
in the 0.72Mn(H) alloy (see Figure 1b,c and Figure 2). To quantitatively compare the dispersoids,
TEM investigations of the as-extruded Mn-containing alloys were conducted (Figure 2), because the
aging treatment at 180
◦
C had no effect on the coarsening of the dispersoids. The dispersoids in the
matrix were identified as
α
-Al(FeMn)Si [
10
,
11
], and their quantitative results are presented in Table 2.
The average equivalent diameter (
D
) and number density (
Nv
) of the dispersoids in the 0.72Mn(L)
alloy were measured to be 40 nm and 430
µ
m
−3
, respectively (versus
D
of 146 nm and
Nv
of 11
µ
m
−3
in 0.72Mn(H) alloy), while the volume fractions of the dispersoids (
fd
) in both alloys were similar
(approximately 0.9%).
Based on the EBSD results of the as-extruded materials, it could be observed that the 0Mn alloy
experienced full recrystallization during extrusion because the grains were coarse, equiaxed, and free
of the substructure (Figure 3a) [
25
]. A <001>recrystallization texture highly dominated in this alloy
(Figure 3a). The average equivalent diameter of the grains in 0Mn alloy were measured as ~220
µ
m.
By contrast, the grain structures of the 0.72Mn(H) and 0.72Mn(L) alloys exhibited deformed and fibrous
grains elongated along the extrusion direction (Figure 3b,c). The deformed grains appeared in one of
the two <111>and <001>directions parallel to the extrusion axis, thereby indicating the evolution
of <111>and <001>fiber textures during the axisymmetric deformation [
26
,
27
]. Evidently, only a
dynamic recovery occurred in the 0.72Mn(H) and 0.72Mn(L) alloys during the extrusion, and the
0.72Mn(L) alloy exhibited a lower recovery level and a higher fraction of substructures (misorientation
angle ranges between 5◦and 15◦) than did the 0.72Mn(H) alloy.

Materials 2020,13, 5295 5 of 13
Materials 2020, 13, x FOR PEER REVIEW 4 of 13
because the aging treatment at 180 °C had no effect on the coarsening of the dispersoids. The
dispersoids in the matrix were identified as α-Al(FeMn)Si [10,11], and their quantitative results are
presented in Table 2. The average equivalent diameter (
) and number density () of the dispersoids
in the 0.72Mn(L) alloy were measured to be 40 nm and 430 µm
−3
, respectively
(versus 
 of 146 nm
and  of 11 µm
−3
in 0.72Mn(H) alloy), while the volume fractions of the dispersoids (ʄ) in both
alloys were similar (approximately 0.9%).
Figure 1. Bright-field TEM images of precipitates and dispersoids in aluminum matrices of T5-treated
(a) 0Mn, (b) 0.72Mn(H), and (c) 0.72Mn(L) alloys.
Figure 2. TEM images showing α-Al(FeMn)Si dispersoids embedded in aluminum matrices of (a)
0.72Mn(H) and (b) 0.72Mn(L) alloys. Images were taken near [001]
Al
zone axes.
Table 2. Quantitative TEM results of β” precipitates and α-Al(FeMn)Si dispersoids in three alloys.
Alloy β” phase
α
-Al(FeMn)Si Dispersoids
, nm
3
, nm
−3
,% 
 , µm
−3
,%
0Mn 410 1.26 × 10
–5
0.52 - - -
0.72Mn(H) 459 1.40 × 10
–5
0.62 146 11 0.91
0.72Mn(L) 468 1.33 × 10
–5
0.64 40 430 0.88
Based on the EBSD results of the as-extruded materials, it could be observed that the 0Mn alloy
experienced full recrystallization during extrusion because the grains were coarse, equiaxed, and free
of the substructure (Figure 3a) [25]. A <001> recrystallization texture highly dominated in this alloy
(Figure 3a). The average equivalent diameter of the grains in 0Mn alloy were measured as ~220 µm.
Figure 2.
TEM images showing
α
-Al(FeMn)Si dispersoids embedded in aluminum matrices of (
a
)
0.72Mn(H) and (b) 0.72Mn(L) alloys. Images were taken near [001]Al zone axes.
Materials 2020, 13, x FOR PEER REVIEW 5 of 13
By contrast, the grain structures of the 0.72Mn(H) and 0.72Mn(L) alloys exhibited deformed and
fibrous grains elongated along the extrusion direction (Figure 3b,c). The deformed grains appeared
in one of the two <111> and <001> directions parallel to the extrusion axis, thereby indicating the
evolution of <111> and <001> fiber textures during the axisymmetric deformation [26,27]. Evidently,
only a dynamic recovery occurred in the 0.72Mn(H) and 0.72Mn(L) alloys during the extrusion, and
the 0.72Mn(L) alloy exhibited a lower recovery level and a higher fraction of substructures
(misorientation angle ranges between 5° and 15°) than did the 0.72Mn(H) alloy.
Figure 3. EBSD (inverse pole figure) maps representing extruded grain structures: (a) 0Mn, (b)
0.72Mn(H), and (c) 0.72Mn(L) alloys.
Note that the EBSD scan area for 0Mn is much larger than
for 0.72Mn(H) and 0.72Mn(L) alloys.
Figure 4a displays the compressive true stress-strain curves at room temperature of the alloys in
T5 state. Three alloys present similar work hardening behavior at the beginning of the compression
test. With increasing the strain, the stress of the 0.72Mn(L) alloy was moderately higher than that of
the 0.72Mn(H) alloy, but the stresses of both alloys were remarkably higher than that of the 0Mn base
alloy. The obtained YSs (determined at 0.2% offset strain) of three alloys are shown in Figure 4b. It is
evident that both Mn-containing alloys (0.72Mn(H) and 0.72Mn(L)) exhibited higher YS than the 0Mn
base alloy (310 MPa versus 270 MPa, respectively). The YS of 6082 alloys under T5 conditions was
mainly controlled by the nanoscale β” precipitates [22]. The higher strength of the Mn-containing
alloys was primarily attributed to the higher number density and larger volume fraction of β”
precipitates in the matrix (Table 2). Probably, the enhanced diffusion of solutes and accelerated
nucleation and growth of β” precipitates in the deformed grains of both Mn-containing alloys led to
the precipitation of a larger fraction of the β” phase than in the recrystallized 0Mn alloy during the
subsequent aging treatment [9]. In addition, the higher strength of the Mn-containing alloys was
partially contributed by their fibrous grain structure and texture compared to the fully recrystallized
grains of the 0Mn alloy (Figure 3). Owing to the relatively large size and low density of dispersoids
in the Mn-containing alloys, the strengthening effect of the dispersoids may have had a minor effect
on the room temperature strength of the alloys relative to the predominant effect of β” precipitates.
The fact that both 0.72Mn(L) and 0.72Mn(H) alloys under T5 condition exhibited similar YS indicated
Figure 3.
EBSD (inverse pole figure) maps representing extruded grain structures: (
a
) 0Mn,
(
b
) 0.72Mn(H), and (
c
) 0.72Mn(L) alloys. Note that the EBSD scan area for 0Mn is much larger
than for 0.72Mn(H) and 0.72Mn(L) alloys.
Figure 4a displays the compressive true stress-strain curves at room temperature of the alloys in
T5 state. Three alloys present similar work hardening behavior at the beginning of the compression
test. With increasing the strain, the stress of the 0.72Mn(L) alloy was moderately higher than that of the
0.72Mn(H) alloy, but the stresses of both alloys were remarkably higher than that of the 0Mn base alloy.
The obtained YSs (determined at 0.2% offset strain) of three alloys are shown in Figure 4b. It is evident
that both Mn-containing alloys (0.72Mn(H) and 0.72Mn(L)) exhibited higher YS than the 0Mn base
alloy (310 MPa versus 270 MPa, respectively). The YS of 6082 alloys under T5 conditions was mainly

Materials 2020,13, 5295 6 of 13
controlled by the nanoscale
β
” precipitates [
22
]. The higher strength of the Mn-containing alloys
was primarily attributed to the higher number density and larger volume fraction of
β
” precipitates
in the matrix (Table 2). Probably, the enhanced diffusion of solutes and accelerated nucleation and
growth of
β
” precipitates in the deformed grains of both Mn-containing alloys led to the precipitation
of a larger fraction of the
β
” phase than in the recrystallized 0Mn alloy during the subsequent aging
treatment [
9
]. In addition, the higher strength of the Mn-containing alloys was partially contributed by
their fibrous grain structure and texture compared to the fully recrystallized grains of the 0Mn alloy
(Figure 3). Owing to the relatively large size and low density of dispersoids in the Mn-containing
alloys, the strengthening effect of the dispersoids may have had a minor effect on the room temperature
strength of the alloys relative to the predominant effect of
β
” precipitates. The fact that both 0.72Mn(L)
and 0.72Mn(H) alloys under T5 condition exhibited similar YS indicated the limited strengthening
effect of the dispersoids in the presence of a high number density of β” precipitates.
Materials 2020, 13, x FOR PEER REVIEW 6 of 13
the limited strengthening effect of the dispersoids in the presence of a high number density of β"
precipitates.
Figure 4. (a) Compressive true stress-strain curves and (b) compressive YS at room
temperature of the alloys under T5 conditions.
3.2. Instantaneous High-Temperature Strength during Thermal Exposure
The typical compressive true stress-strain curves at 400 °C are exemplarily displayed in Figure
5a. Obviously, the alloys show almost no work hardening behaviour at high temperature. Figure 5b,c
shows the instantaneous YS of the three alloys at various temperatures and thermal exposures for 0.5
and 2 h. For both thermal exposure times, the strength of the alloys decreased with increasing
exposure temperature. At 200 °C, all alloys thermally exposed for 0.5 h displayed a similar YS of
approximately 220 MPa (Figure 5b). With increasing exposure time to 2 h, the YS of the 0Mn alloy
increased slightly to 225 MPa, while the Mn-containing alloys displayed a slight decrease in their
strength, and the 0.72Mn(L) alloy showed the lowest strength (202 MPa) (Figure 5c).
Compared to the strengths determined at 200 °C, the strengths of all three alloys at 300 °C
decreased significantly. The YS at 300 °C and the thermal exposure of 0.5 h varied between 70 and 80
MPa. With increasing exposure time to 2 h, the alloys experienced a further reduction in their YSs.
Interestingly, the 0.72Mn(L) alloy began to show its advantage in terms of strength, and its YS of 70
MPa for a 2 h exposure was higher than those of the other two alloys. The 0.72Mn(H) alloy exhibited
the sharpest decrease in its YS (48 MPa). In addition, the YS of the 0Mn alloy reached 64 MPa.
With further increasing temperature up to 450 °C, the strength of all alloys continued to decrease.
The instantaneous YS of the 0Mn alloy decreased sharply at such high temperatures and attained
only 11 MPa after an exposure at 450 °C; this represented only 5% of the YS after the exposure at 200
°C. The 0.72Mn(L) alloy showed the highest YS at these high temperatures; its YS values at each
temperature for both exposure times were similar. Specifically, the YS values of the 0.72Mn(L) alloy
after 2 h exposure reached 52, 40, and 31 MPa at 350, 400, and 450 °C, respectively, which were much
higher than those of the 0.72Mn(H) and 0Mn alloys at the corresponding temperatures. For instance,
after being exposed to 400–450 °C, the YS values of the 0.72Mn(L) alloy were 60%–75% and 150%–
180% higher than those of the 0.72Mn(H) and 0Mn alloys, respectively. This demonstrates that the
addition of Mn and the formation of thermally stable dispersoids had a great impact on the
instantaneous YS at high temperatures (350–450 °C).
Figure 4.
(
a
) Compressive true stress-strain curves and (
b
) compressive YS at room temperature of the
alloys under T5 conditions.
3.2. Instantaneous High-Temperature Strength during Thermal Exposure
The typical compressive true stress-strain curves at 400
◦
C are exemplarily displayed in Figure 5a.
Obviously, the alloys show almost no work hardening behaviour at high temperature. Figure 5b,c
shows the instantaneous YS of the three alloys at various temperatures and thermal exposures for
0.5 and 2 h. For both thermal exposure times, the strength of the alloys decreased with increasing
exposure temperature. At 200
◦
C, all alloys thermally exposed for 0.5 h displayed a similar YS of
approximately 220 MPa (Figure 5b). With increasing exposure time to 2 h, the YS of the 0Mn alloy
increased slightly to 225 MPa, while the Mn-containing alloys displayed a slight decrease in their
strength, and the 0.72Mn(L) alloy showed the lowest strength (202 MPa) (Figure 5c).
Compared to the strengths determined at 200
◦
C, the strengths of all three alloys at 300
◦
C
decreased significantly. The YS at 300
◦
C and the thermal exposure of 0.5 h varied between 70 and
80 MPa. With increasing exposure time to 2 h, the alloys experienced a further reduction in their
YSs. Interestingly, the 0.72Mn(L) alloy began to show its advantage in terms of strength, and its YS of
70 MPa for a 2 h exposure was higher than those of the other two alloys. The 0.72Mn(H) alloy exhibited
the sharpest decrease in its YS (48 MPa). In addition, the YS of the 0Mn alloy reached 64 MPa.
With further increasing temperature up to 450
◦
C, the strength of all alloys continued to decrease.
The instantaneous YS of the 0Mn alloy decreased sharply at such high temperatures and attained
only 11 MPa after an exposure at 450
◦
C; this represented only 5% of the YS after the exposure at
200
◦
C. The 0.72Mn(L) alloy showed the highest YS at these high temperatures; its YS values at each
temperature for both exposure times were similar. Specifically, the YS values of the 0.72Mn(L) alloy
after 2 h exposure reached 52, 40, and 31 MPa at 350, 400, and 450 ◦C, respectively, which were much
higher than those of the 0.72Mn(H) and 0Mn alloys at the corresponding temperatures. For instance,
after being exposed to 400–450
◦
C, the YS values of the 0.72Mn(L) alloy were 60%–75% and 150%–180%

Materials 2020,13, 5295 7 of 13
higher than those of the 0.72Mn(H) and 0Mn alloys, respectively. This demonstrates that the addition
of Mn and the formation of thermally stable dispersoids had a great impact on the instantaneous YS at
high temperatures (350–450 ◦C).
Materials 2020, 13, x FOR PEER REVIEW 7 of 13
Figure 5. (a) Typical compressive true stress-strain curves at 400 °C and (b,c) instantaneous YS
of three alloys at various temperatures, measured after being exposed to the test temperature for
(b) 0.5 h and (c) 2 h.
3.3. Residual Room-Temperature Strength after Thermal Exposure
The residual mechanical properties of the aluminum structure after the fire are a key parameter
for evaluating the stability and integrity of the structure as well as the possibility of repair,
reinforcement, and reuse [2,21]. To characterize the residual mechanical response, compressive YS
tests were conducted on the alloys at room temperature after their exposure to elevated temperatures
for two exposure times (0.5 and 2 h). The results are shown in Figure 6.
After the exposure at 200 °C, the three experimental alloys still exhibited high residual YSs,
ranging between 280 and 300 MPa, which were at a similar level as the YSs under T5 conditions
(Figure 4). Increasing the exposure time to 2 h exerted a negligible effect on the YSs of the alloys
(Figure 6b). These results implied that the exposure to 200 °C for the investigated exposure times had
no apparent deleterious effect on the mechanical properties of the aluminum structures. Therefore,
they could be reused after fire.
Consistent with the high-temperature YS results, the exposure of the alloys to 300 °C resulted in
a significant reduction in the residual YS (Figure 6). After thermal exposure for 2 h, the 0.72Mn(L)
alloy reached a YS of 150 MPa, which largely exceeded those of the 0Mn and 0.72Mn(H) alloys
(residual YSs of 105 and 95 MPa, respectively). It was evident that the residual YSs of the three alloys
after the exposure to 300 °C were only 35%–45% of the original YSs under T5 conditions.
At higher exposure temperatures (350–450 °C), the residual YS of the 0.72Mn(L) alloy remained
basically constant at the level of 115–120 MPa regardless of the exposure temperature and time. By
contrast, the 0.72Mn(H) alloy exposed to 350 °C for 0.5 h exhibited a residual YS of 95 MPa.
Furthermore, with increasing exposure temperature to 400–450 °C and exposure time to 2 h, the
residual YS of this alloy decreased until it stabilized at 80 MPa. The residual mechanical response of
Figure 5.
(
a
) Typical compressive true stress-strain curves at 400
◦
C and (
b
,
c
) instantaneous YS of three
alloys at various temperatures, measured after being exposed to the test temperature for (
b
) 0.5 h and
(c) 2 h.
3.3. Residual Room-Temperature Strength after Thermal Exposure
The residual mechanical properties of the aluminum structure after the fire are a key parameter for
evaluating the stability and integrity of the structure as well as the possibility of repair, reinforcement,
and reuse [
2
,
21
]. To characterize the residual mechanical response, compressive YS tests were conducted
on the alloys at room temperature after their exposure to elevated temperatures for two exposure times
(0.5 and 2 h). The results are shown in Figure 6.
After the exposure at 200
◦
C, the three experimental alloys still exhibited high residual YSs,
ranging between 280 and 300 MPa, which were at a similar level as the YSs under T5 conditions
(Figure 4). Increasing the exposure time to 2 h exerted a negligible effect on the YSs of the alloys
(Figure 6b). These results implied that the exposure to 200
◦
C for the investigated exposure times had
no apparent deleterious effect on the mechanical properties of the aluminum structures. Therefore,
they could be reused after fire.
Consistent with the high-temperature YS results, the exposure of the alloys to 300
◦
C resulted in a
significant reduction in the residual YS (Figure 6). After thermal exposure for 2 h, the 0.72Mn(L) alloy
reached a YS of 150 MPa, which largely exceeded those of the 0Mn and 0.72Mn(H) alloys (residual
YSs of 105 and 95 MPa, respectively). It was evident that the residual YSs of the three alloys after the
exposure to 300 ◦C were only 35%–45% of the original YSs under T5 conditions.

Materials 2020,13, 5295 8 of 13
At higher exposure temperatures (350–450
◦
C), the residual YS of the 0.72Mn(L) alloy remained
basically constant at the level of 115–120 MPa regardless of the exposure temperature and time.
By contrast, the 0.72Mn(H) alloy exposed to 350
◦
C for 0.5 h exhibited a residual YS of 95 MPa.
Furthermore, with increasing exposure temperature to 400–450
◦
C and exposure time to 2 h, the residual
YS of this alloy decreased until it stabilized at 80 MPa. The residual mechanical response of the 0Mn
alloy showed the worst performance. For instance, the 0Mn alloy exposed to 350
◦
C for 2 h attained a
residual YS of 68.5 MPa; however, following the exposure to 400–450
◦
C, the residual YS decreased
to 49 MPa, which represented only 18.5% of the original YS obtained prior to the high-temperature
exposure. Hence, the 0.72Mn(L) alloy is a promising candidate for fire-susceptible aluminum structures,
owing to its 50% and 140% higher residual YSs after exposures to 400–450
◦
C for 2 h than those of the
0.72Mn(H) and 0Mn alloys, respectively.
Materials 2020, 13, x FOR PEER REVIEW 8 of 13
the 0Mn alloy showed the worst performance. For instance, the 0Mn alloy exposed to 350 °C for 2 h
attained a residual YS of 68.5 MPa; however, following the exposure to 400–450 °C, the residual YS
decreased to 49 MPa, which represented only 18.5% of the original YS obtained prior to the high-
temperature exposure. Hence, the 0.72Mn(L) alloy is a promising candidate for fire-susceptible
aluminum structures, owing to its 50% and 140% higher residual YSs after exposures to 400–450 °C
for 2 h than those of the 0.72Mn(H) and 0Mn alloys, respectively.
Figure 6. Residual room temperature YS of the experimental alloys; all alloys were pre-exposed to the
test temperature for (a) 0.5 h and (b) 2 h.
4. Discussion
Typically, Al–Mg–Si 6082 alloys experience predominant strengthening by the precipitation of
nanoscale precipitates (particularly β”), which are formed during aging treatment [21]. Therefore, the
precipitate stability during high temperature exposure determines the capacity of the aluminum
structures to bear a certain amount of load, i.e., its strength during and after a fire [1]. With the
addition of an adequate Mn to 6082 alloys and the appropriate heat treatment, a large number of
thermally stable α-Al(FeMn)Si dispersoids can be formed during homogenization [11,20]; these
contribute to preserving the high-temperature strength through dispersoid strengthening [10,11] and
delayed restoration processes [25]. The extent of microstructural and mechanical stability of the
aluminum alloys during high temperature exposure depends on the temperature and time [1].
At 200 °C, the overall instantaneous and residual YSs of all three alloys are still high and
comparable, indicating that the β” precipitates formed during aging (T5 condition) control the
strength of the alloys. In other words, the mechanical properties of the three alloys after exposures to
200 °C are not markedly deteriorated yet. The 0Mn alloy exhibits a slight increase in the instantaneous
YS with prolonged thermal exposure, unlike the Mn-containing alloys (Figure 5). This can be
explained by considering the precipitation structures under T5 conditions. Owing to the slower
precipitation rate of the β” phase in the 0Mn alloy with recrystallized grain structure than that in the
Mn-containing alloys with deformed grain structures, less dense and relatively finer β” precipitates
were observed in the former alloy (Table 2). The prolonged exposure time for 2 h at 200 °C causes a
further growth of the β” precipitates in the 0Mn alloy, thereby slightly increasing the YS relative to
the value of the 0.5 h exposure. However, the Mn-containing alloys exhibit the over-aging
phenomenon after a longer exposure. The higher amount of energy stored in the Mn-containing
alloys (0.72Mn(L)) in the form of dislocations and subgrain boundaries (Figure 3) accelerates the
coarsening of β” precipitates, thereby resulting in a moderate decrease in the strength and lower YS
compared to that of the Mn-free 0Mn alloy.
The three alloys exposed to 300 °C exhibit significant strength reductions, which indicates
relevant changes in the strengthening precipitates. In fact, the TEM results of the alloys exposed to
300 °C for 2 h reveal the presence of coarse rod-shaped precipitates (Figure 7), which were identified
as β’ phase based on their size and morphology [24]. The coarse β’ precipitates are known to be less
effective in alloy strengthening [24,28,29]. Moreover, by comparing the characteristics of the β’
precipitates of the various alloys, it can be observed that the β’ precipitates in the 0Mn alloy are finer
Figure 6.
Residual room temperature YS of the experimental alloys; all alloys were pre-exposed to the
test temperature for (a) 0.5 h and (b) 2 h.
4. Discussion
Typically, Al–Mg–Si 6082 alloys experience predominant strengthening by the precipitation of
nanoscale precipitates (particularly
β
”), which are formed during aging treatment [
21
]. Therefore,
the precipitate stability during high temperature exposure determines the capacity of the aluminum
structures to bear a certain amount of load, i.e., its strength during and after a fire [
1
]. With the addition
of an adequate Mn to 6082 alloys and the appropriate heat treatment, a large number of thermally
stable
α
-Al(FeMn)Si dispersoids can be formed during homogenization [
11
,
20
]; these contribute
to preserving the high-temperature strength through dispersoid strengthening [
10
,
11
] and delayed
restoration processes [
25
]. The extent of microstructural and mechanical stability of the aluminum
alloys during high temperature exposure depends on the temperature and time [1].
At 200
◦
C, the overall instantaneous and residual YSs of all three alloys are still high and
comparable, indicating that the
β
” precipitates formed during aging (T5 condition) control the strength
of the alloys. In other words, the mechanical properties of the three alloys after exposures to 200
◦
C are
not markedly deteriorated yet. The 0Mn alloy exhibits a slight increase in the instantaneous YS with
prolonged thermal exposure, unlike the Mn-containing alloys (Figure 5). This can be explained by
considering the precipitation structures under T5 conditions. Owing to the slower precipitation rate
of the
β
” phase in the 0Mn alloy with recrystallized grain structure than that in the Mn-containing
alloys with deformed grain structures, less dense and relatively finer
β
” precipitates were observed
in the former alloy (Table 2). The prolonged exposure time for 2 h at 200
◦
C causes a further growth
of the
β
” precipitates in the 0Mn alloy, thereby slightly increasing the YS relative to the value of
the 0.5 h exposure. However, the Mn-containing alloys exhibit the over-aging phenomenon after a
longer exposure. The higher amount of energy stored in the Mn-containing alloys (0.72Mn(L)) in the
form of dislocations and subgrain boundaries (Figure 3) accelerates the coarsening of
β
” precipitates,

Materials 2020,13, 5295 9 of 13
thereby resulting in a moderate decrease in the strength and lower YS compared to that of the Mn-free
0Mn alloy.
The three alloys exposed to 300
◦
C exhibit significant strength reductions, which indicates relevant
changes in the strengthening precipitates. In fact, the TEM results of the alloys exposed to 300
◦
C for
2 h reveal the presence of coarse rod-shaped precipitates (Figure 7), which were identified as
β
’ phase
based on their size and morphology [
24
]. The coarse
β
’ precipitates are known to be less effective in
alloy strengthening [
24
,
28
,
29
]. Moreover, by comparing the characteristics of the
β
’ precipitates of the
various alloys, it can be observed that the
β
’ precipitates in the 0Mn alloy are finer and denser than those
in the Mn-containing alloys (Figure 7). The difference in the sizes and densities of the
β
’ precipitates
between the 0Mn and Mn-containing alloys is likely associated with their initial as-extruded grain
structures and, more specifically, with the higher rate of solute diffusion in the deformed grains of
the Mn-containing alloys, which results in a quick coarsening of the
β
’ precipitates during thermal
exposure [
9
]. This also explains why the 0Mn alloy possesses a higher YS than the 0.72Mn(H) alloy (for
instance, 65 MPa versus 47 MPa after 2 h exposure at 300
◦
C, respectively). The
β
’ precipitates in the
0.72Mn(L) alloy are the same as those in the 0.72Mn(H) alloy. However, the 0.72Mn(L) alloy possesses
a high number of fine
α
-dispersoids in the matrix (Figure 7c). It is evident that the dispersoids show
no coarsening tendency during the thermal exposure and thus provide additional strength. Therefore,
this alloy exhibits the highest instantaneous YS after a 2 h exposure (70 MPa; Figure 5). Moreover,
the residual YS at 300
◦
C of this alloy is also substantially higher than those of the other two alloys
(Figure 6) owing to the complementary strengthening effect of the dispersoids.
Materials 2020, 13, x FOR PEER REVIEW 9 of 13
and denser than those in the Mn-containing alloys (Figure 7). The difference in the sizes and densities
of the β’ precipitates between the 0Mn and Mn-containing alloys is likely associated with their initial
as-extruded grain structures and, more specifically, with the higher rate of solute diffusion in the
deformed grains of the Mn-containing alloys, which results in a quick coarsening of the β’ precipitates
during thermal exposure [9]. This also explains why the 0Mn alloy possesses a higher YS than the
0.72Mn(H) alloy (for instance, 65 MPa versus 47 MPa after 2 h exposure at 300 °C, respectively). The
β’ precipitates in the 0.72Mn(L) alloy are the same as those in the 0.72Mn(H) alloy. However, the
0.72Mn(L) alloy possesses a high number of fine α-dispersoids in the matrix (Figure 7c). It is evident
that the dispersoids show no coarsening tendency during the thermal exposure and thus provide
additional strength. Therefore, this alloy exhibits the highest instantaneous YS after a 2 h exposure
(70 MPa; Figure 5). Moreover, the residual YS at 300 °C of this alloy is also substantially higher than
those of the other two alloys (Figure 6) owing to the complementary strengthening effect of the
dispersoids.
At higher exposure temperatures (350–450 °C), the microstructures of the alloys keep changing
with decreasing strength of the alloys. Figure 8 shows the typical TEM microstructure of the
0.72Mn(L) alloy exposed to 400 °C for 2 h. The Mg
2
Si precursor precipitates experience significant
changes; they are transformed into µm-sized equilibrium β-Mg
2
Si particles (Figure 8a), while the α-
dispersoids show no apparent coarsening (Figure 8b), because their dimensions are comparable to
those in the as-extruded materials (Figure 2). The relatively high diffusion rates of Mg and Si in the
aluminum matrix at such high temperatures cause the transformation of Mg
2
Si precursor precipitates
into coarse equilibrium β particles, which lost their major strengthening effect [24]. In this case, the
presence of thermally stable α-dispersoids contributes effectively to the strength of the Mn-
containing alloys, becoming the dominant strengthening mechanism. Therefore, the dispersoid-free
0Mn alloy has very low instantaneous and residual YSs at elevated temperatures, while the 0.72Mn(L)
alloy with fine and dense dispersoids has superior instantaneous and residual YSs, which results in
the best mechanical response among the three alloys.
Figure 7. Bright-field TEM images of (a) 0Mn, (b) 0.72Mn(H), and (c) 0.72Mn(L) alloys exposed to 300
°C for 2 h after T5 treatment.
Figure 7.
Bright-field TEM images of (
a
) 0Mn, (
b
) 0.72Mn(H), and (
c
) 0.72Mn(L) alloys exposed to
300 ◦C for 2 h after T5 treatment.
At higher exposure temperatures (350–450
◦
C), the microstructures of the alloys keep changing
with decreasing strength of the alloys. Figure 8shows the typical TEM microstructure of the 0.72Mn(L)
alloy exposed to 400
◦
C for 2 h. The Mg
2
Si precursor precipitates experience significant changes; they
are transformed into
µ
m-sized equilibrium
β
-Mg
2
Si particles (Figure 8a), while the
α
-dispersoids
show no apparent coarsening (Figure 8b), because their dimensions are comparable to those in the
as-extruded materials (Figure 2). The relatively high diffusion rates of Mg and Si in the aluminum
matrix at such high temperatures cause the transformation of Mg
2
Si precursor precipitates into coarse
equilibrium
β
particles, which lost their major strengthening effect [
24
]. In this case, the presence
of thermally stable
α
-dispersoids contributes effectively to the strength of the Mn-containing alloys,
becoming the dominant strengthening mechanism. Therefore, the dispersoid-free 0Mn alloy has very
low instantaneous and residual YSs at elevated temperatures, while the 0.72Mn(L) alloy with fine and
dense dispersoids has superior instantaneous and residual YSs, which results in the best mechanical
response among the three alloys.

Materials 2020,13, 5295 10 of 13
Materials 2020, 13, x FOR PEER REVIEW 10 of 13
Figure 8. Bright-field TEM images showing (a) coarsened β-Mg
2
Si particles and (b) α-Al(FeMn)Si
dispersoids in 0.72Mn(L) alloy exposed to 400 °C for 2 h after T5 treatment. Images were taken near
[001]
Al
zone axes.
The grain structures of the three alloys after the high-temperature exposure (400 °C for 2 h) were
investigated with EBSD (Figure 9). The 0Mn alloy maintains fully recrystallized grain structure
(Figure 9a) because recrystallization has already occurred during extrusion (Figure 3a). The average
equivalent diameter of grains in 0Mn alloy is ~220 µm, which is quite the same as that of the as-
extruded 0Mn alloy (Figure 3a). This implies that no appreciable grain growth occurred during the
high-temperature exposure. The 0.72Mn(H) alloy retains a large part of the recovered grain structure
with some recrystallized grains on the grain boundaries (see arrows in Figure 9b). As previously
mentioned, the 0.72Mn(H) alloy contains a number of α-dispersoids, which promote the
recrystallization retardation owing to their pinning effect; nevertheless, it is less effective than that of
the 0.72Mn(L) alloy. The fact that the 0.72Mn(H) alloy exhibits a higher residual YS than the 0Mn
alloy at higher exposure temperatures (350–450 °C; Figure 6) can be partially attributed to their
recovered grain structures, which enable better subgrain and strain hardening [9]. On the other hand,
the 0.72Mn(L) alloy maintains its deformed and fibrous grain structure (Figure 9c) owing to the
strong pinning effect of a large number of α-dispersoids, which inhibit recrystallization during the
high-temperature exposure. The stable and high residual YSs at these high temperatures are
contributed to a certain extent by the non-recrystallized grain structure owing to the subgrain
strengthening effect.
In brief, at a low exposure temperature (200 °C), the fine nanoscale β” precipitates control the
strength of Al–Mg–Si 6082 alloys, and the mechanical properties of aluminum structural alloys
remain almost undeteriorated. With increasing exposure temperature to 300 °C, the β” precipitates
transformed into coarse β’ precipitates, thereby resulting in a significant reduction in the alloy
strength. Owing to a large number of fine and dense dispersoids, the 0.72Mn(L) alloy begins to show
its superior mechanical response and exhibits remarkably higher instantaneous and residual YSs than
the base 0Mn alloy and 0.72Mn(H) alloy, owing to the complementary strengthening effect of the
thermally stable dispersoids. At higher exposure temperatures (350–450 °C), the Mg
2
Si precursor
precipitates transform into coarse equilibrium β-Mg
2
Si particles and lose their strengthening effect.
Moreover, the strength of the base 0Mn alloy decreases with increasing exposure temperature, and
the alloy possesses very low instantaneous and residual YSs at high temperatures. On the other hand,
the dispersoids show no apparent coarsening and effectively contribute to alloy strength. Therefore,
the 0.72Mn(L) alloy with fine and dense dispersoids exhibits the best mechanical response at 350–450
°C among the three alloys. Hence, introducing thermally stable dispersoids by appropriate Mn
alloying and heat treatments can be a cost-effective and promising approach for improving the
mechanical response of aluminum structural alloys during high temperature exposure.
Figure 8.
Bright-field TEM images showing (
a
) coarsened
β
-Mg
2
Si particles and (
b
)
α
-Al(FeMn)Si
dispersoids in 0.72Mn(L) alloy exposed to 400
◦
C for 2 h after T5 treatment. Images were taken near
[001]Al zone axes.
The grain structures of the three alloys after the high-temperature exposure (400
◦
C for 2 h)
were investigated with EBSD (Figure 9). The 0Mn alloy maintains fully recrystallized grain structure
(Figure 9a) because recrystallization has already occurred during extrusion (Figure 3a). The average
equivalent diameter of grains in 0Mn alloy is ~220
µ
m, which is quite the same as that of the
as-extruded 0Mn alloy (Figure 3a). This implies that no appreciable grain growth occurred during the
high-temperature exposure. The 0.72Mn(H) alloy retains a large part of the recovered grain structure
with some recrystallized grains on the grain boundaries (see arrows in Figure 9b). As previously
mentioned, the 0.72Mn(H) alloy contains a number of
α
-dispersoids, which promote the recrystallization
retardation owing to their pinning effect; nevertheless, it is less effective than that of the 0.72Mn(L)
alloy. The fact that the 0.72Mn(H) alloy exhibits a higher residual YS than the 0Mn alloy at higher
exposure temperatures (350–450
◦
C; Figure 6) can be partially attributed to their recovered grain
structures, which enable better subgrain and strain hardening [
9
]. On the other hand, the 0.72Mn(L)
alloy maintains its deformed and fibrous grain structure (Figure 9c) owing to the strong pinning
effect of a large number of
α
-dispersoids, which inhibit recrystallization during the high-temperature
exposure. The stable and high residual YSs at these high temperatures are contributed to a certain
extent by the non-recrystallized grain structure owing to the subgrain strengthening effect.
In brief, at a low exposure temperature (200
◦
C), the fine nanoscale
β
” precipitates control
the strength of Al–Mg–Si 6082 alloys, and the mechanical properties of aluminum structural alloys
remain almost undeteriorated. With increasing exposure temperature to 300
◦
C, the
β
” precipitates
transformed into coarse
β
’ precipitates, thereby resulting in a significant reduction in the alloy strength.
Owing to a large number of fine and dense dispersoids, the 0.72Mn(L) alloy begins to show its superior
mechanical response and exhibits remarkably higher instantaneous and residual YSs than the base 0Mn
alloy and 0.72Mn(H) alloy, owing to the complementary strengthening effect of the thermally stable
dispersoids. At higher exposure temperatures (350–450
◦
C), the Mg
2
Si precursor precipitates transform
into coarse equilibrium
β
-Mg
2
Si particles and lose their strengthening effect. Moreover, the strength of
the base 0Mn alloy decreases with increasing exposure temperature, and the alloy possesses very low
instantaneous and residual YSs at high temperatures. On the other hand, the dispersoids show no
apparent coarsening and effectively contribute to alloy strength. Therefore, the 0.72Mn(L) alloy with
fine and dense dispersoids exhibits the best mechanical response at 350–450
◦
C among the three alloys.
Hence, introducing thermally stable dispersoids by appropriate Mn alloying and heat treatments
can be a cost-effective and promising approach for improving the mechanical response of aluminum
structural alloys during high temperature exposure.

Materials 2020,13, 5295 11 of 13
Materials 2020, 13, x FOR PEER REVIEW 11 of 13
Figure 9. EBSD (inverse pole figure) maps of (a) 0Mn, (b) 0.72Mn(H), and (c) 0.72Mn(L) alloys,
exposed to 400 °C for 2 h. Note that the EBSD scan area for 0Mn is much larger than for 0.72Mn(H)
and 0.72Mn(L) alloys.
5. Conclusions
 The addition of Mn in a typical high-temperature homogenization treatment produced a
number of α-Al(FeMn)Si dispersoids in the 0.72Mn(H) alloy, which improved the mechanical
response during thermal exposure at 350–450 °C relative to that of the Mn-free base alloy.
Moreover, the low-temperature homogenization treatment resulted in a high density of fine
dispersoids in the 0.72Mn(L) alloy, which further improved the mechanical response
substantially.
 At a low exposure temperature (200 °C), the instantaneous and residual YSs of the alloys were
mainly governed by the β” precipitates, and the mechanical properties of the aluminum
structural alloys were not markedly affected compared to their original strength under T5
condition.
 At a high exposure temperature of 300 °C, the β” precipitates transformed into coarse β’
precipitates, thereby resulting in a significant reduction in the alloy strength. The strength
improvement due to the presence of fine and dense dispersoids became quite evident,
because the instantaneous and residual strengths of the corresponding alloy (0.72Mn(L))
were higher than those of the alloys without dispersoids (0Mn) or with coarse dispersoids
(0.72Mn(H)).
 At higher exposure temperatures (350–450 °C), the Mg
2
Si precursor precipitates transformed
into coarse equilibrium β-Mg
2
Si particles and lost their strengthening effect, while the
dispersoids resisted the coarsening and became the dominant strengthening contributor. The
0.72Mn(L) alloy containing fine and dense dispersoids displayed far superior instantaneous
and residual YSs compared to the other two alloy variants and therefore the best mechanical
response during high temperature exposure.
 The presence of thermally stable dispersoids effectively retarded the recrystallization during
high temperature exposure, which improved the high-temperature mechanical properties to
a certain extent.
Figure 9.
EBSD (inverse pole figure) maps of (
a
) 0Mn, (
b
) 0.72Mn(H), and (
c
) 0.72Mn(L) alloys, exposed
to 400
◦
C for 2 h. Note that the EBSD scan area for 0Mn is much larger than for 0.72Mn(H) and
0.72Mn(L) alloys.
5. Conclusions
•
The addition of Mn in a typical high-temperature homogenization treatment produced a number
of
α
-Al(FeMn)Si dispersoids in the 0.72Mn(H) alloy, which improved the mechanical response
during thermal exposure at 350–450
◦
C relative to that of the Mn-free base alloy. Moreover,
the low-temperature homogenization treatment resulted in a high density of fine dispersoids in
the 0.72Mn(L) alloy, which further improved the mechanical response substantially.
•
At a low exposure temperature (200
◦
C), the instantaneous and residual YSs of the alloys were
mainly governed by the
β
” precipitates, and the mechanical properties of the aluminum structural
alloys were not markedly affected compared to their original strength under T5 condition.
•
At a high exposure temperature of 300
◦
C, the
β
” precipitates transformed into coarse
β
’ precipitates,
thereby resulting in a significant reduction in the alloy strength. The strength improvement due
to the presence of fine and dense dispersoids became quite evident, because the instantaneous
and residual strengths of the corresponding alloy (0.72Mn(L)) were higher than those of the alloys
without dispersoids (0Mn) or with coarse dispersoids (0.72Mn(H)).
•
At higher exposure temperatures (350–450
◦
C), the Mg
2
Si precursor precipitates transformed into
coarse equilibrium
β
-Mg
2
Si particles and lost their strengthening effect, while the dispersoids
resisted the coarsening and became the dominant strengthening contributor. The 0.72Mn(L) alloy
containing fine and dense dispersoids displayed far superior instantaneous and residual YSs
compared to the other two alloy variants and therefore the best mechanical response during high
temperature exposure.
•
The presence of thermally stable dispersoids effectively retarded the recrystallization during
high temperature exposure, which improved the high-temperature mechanical properties to a
certain extent.

Materials 2020,13, 5295 12 of 13
Author Contributions:
J.R.: Methodology, Investigation, Formal analysis, Writing—Original Draft; K.L.:
Methodology, Validation, Writing—Review and Editing; P.R.: Methodology, Resources, Writing—Review and
Editing; N.P.: Methodology, Resources, Writing—Review and Editing; X.-G.C.: Methodology, Writing—Review
and Editing, Supervision, Project administration, Funding acquisition. All authors have read and agreed to the
published version of the manuscript.
Funding:
This research was funded by Natural Sciences and Engineering Research Council of Canada (NSERC)
under Grant No. CRDPJ 514651-17 and Rio Tinto Aluminum through the Research Chair in the Metallurgy of
Aluminum Transformation at the University of Quebec at Chicoutimi.
Conflicts of Interest: The authors declare no conflict of interest.
References
1.
Summers, P.T.; Chen, Y.; Rippe, C.M.; Allen, B.; Mouritz, A.P.; Case, S.W.; Lattimer, B.Y. Overview of
aluminum alloy mechanical properties during and after fires. Fire Sci. Rev. 2015,4, 3.
2.
Chen, Z.H.; Lu, J.; Liu, H.B.; Liao, X.W. Experimental investigation on the post-fire mechanical properties of
structural aluminum alloys 6061-T6 and 7075-T73. Thin Wall Struct. 2016,106, 187–200.
3.
Maljaars, J.; Soetens, F.; Katgerman, L. Constitutive model for aluminum alloys exposed to fire conditions.
Met. Mater. Trans. A 2008,39A, 778–789.
4.
Fogle, E.J.; Lattimer, B.Y.; Feih, S.; Kandare, E.; Mouritz, A.P.; Case, S.W. Compression load failure of
aluminum plates due to fire. Eng. Struct. 2012,34, 155–162.
5.
Hu, J.; Zhang, W.; Fu, D.; Teng, J.; Zhang, H. Improvement of the mechanical properties of Al–Mg–Si alloys
with nano-scale precipitates after repetitive continuous extrusion forming and T8 tempering. J. Mater. Res.
Technol. 2019,8, 5950–5960.
6.
Sekhar, A.P.; Mandal, A.B.; Das, D. Mechanical properties and corrosion behavior of artificially aged Al-Mg-Si
alloy. J. Mater. Res. Technol. 2020,9, 1005–1024.
7.
Kaufman, J.G. Fire Resistance of Aluminum and Aluminum Alloys: Measuring the Effects of Fire Exposure on the
Properties of Aluminum Alloys; ASM International: Materials Park, OH, USA, 2016.
8.
European Center for Standardization. Design of Aluminium Structures. Part 1–2: Structural Fire Design;
European Committee for Standardization: Brussels, Belgium, 2007.
9.
Sigli, C.; De Geuser, F.; Deschamps, A.; Lepinoux, J.; Perez, M. Recent advances in the metallurgy of aluminum
alloys. Part II: Age hardening. Comptes Rendus Phys. 2018,19, 688–709.
10.
Gupta, A.K.; Lloyd, D.J.; Court, S.A. Precipitation hardening in Al-Mg-Si alloys with and without excess Si.
Mater. Sci. Eng. A 2001,316, 11–17.
11.
Fang, X.; Song, M.; Li, K.; Du, Y. Precipitation Sequence of an Aged Al-Mg-Si Alloy. J. Min. Met. B
2010
,
46, 171–180.
12.
Gross, C.T.; Isheim, D.; Vaynman, S.; Fine, M.E.; Chung, Y.W. Design and Development of Lightly Alloyed
Ferritic Fire-Resistant Structural Steels. Met. Mater. Trans. A 2019,50A, 209–219.
13.
Liu, K.; Ma, H.; Chen, X.G. Enhanced elevated-temperature properties via Mo addition in Al-Mn-Mg 3004
alloy. J. Alloys Compd. 2017,694, 354–365.
14.
Liu, K.; Chen, X.G. Evolution of Intermetallics, Dispersoids, and Elevated Temperature Properties at Various
Fe Contents in Al-Mn-Mg 3004 Alloys. Met. Mater. Trans. B 2016,47, 3291–3300.
15.
Liu, K.; Chen, X.G. Development of Al–Mn–Mg 3004 alloy for applications at elevated temperature via
dispersoid strengthening. Mater. Des. 2015,84, 340–350.
16.
Li, Z.; Zhang, Z.; Chen, X.G. Microstructure, elevated-temperature mechanical properties and creep resistance
of dispersoid-strengthened Al-Mn-Mg 3xxx alloys with varying Mg and Si contents. Mater. Sci. Eng. A
2017
,
708, 383–394.
17.
Rakhmonov, J.; Timelli, G.; Fabrizi, A.; Bonollo, F. Effect of V and Zr microalloying, and heat treatment on
microstructure and mechanical properties of secondary Al-7Si-3Cu-0.3Mg alloy. Int. J. Mater. Res.
2018
,
109, 1099–1112.
18.
Du, Q.; Poole, W.J.; Wells, M.A.; Parson, N.C. Microstructure evolution during homogenization of Al-Mn-Fe-Si
alloys: Modeling and experimental results. Acta Mater. 2013,61, 4961–4973.
19.
Li, Y.J.; Arnberg, L. Quantitative study on the precipitation behavior of dispersoids in DC-cast AA3003 alloy
during heating and homogenization. Acta Mater. 2003,51, 3415–3428.

Materials 2020,13, 5295 13 of 13
20.
Li, Y.J.; Muggerud, A.M.F.; Olsen, A.; Furu, T. Precipitation of partially coherent
α
-Al(Mn,Fe)Si dispersoids
and their strengthening effect in AA 3003 alloy. Acta Mater. 2012,60, 1004–1014.
21.
Liu, Y.; Liu, H.B.; Chen, Z.H. Post-fire mechanical properties of aluminum alloy 6082-T6. Constr. Build. Mater.
2019,196, 256–266.
22.
Engler, O.; Marioara, C.D.; Aruga, Y.; Kozuka, M.; Myhr, O.R. Effect of natural ageing or pre-ageing on the
evolution of precipitate structure and strength during age hardening of Al–Mg–Si alloy AA 6016. Mater. Sci.
Eng. A 2019,759, 520–529.
23.
Qian, F.; Jin, S.B.; Sha, G.; Li, Y.J. Enhanced dispersoid precipitation and dispersion strengthening in an Al
alloy by microalloying with Cd. Acta Mater. 2018,157, 114–125.
24.
Edwards, G.A.; Stiller, K.; Dunlop, G.L.; Couper, M.J. The precipitation sequence in Al-Mg-Si alloys. Acta
Mater. 1998,46, 3893–3904.
25.
Mahmoodkhani, Y.; Wells, M.A.; Poole, W.J.; Grajales, L.; Parson, N. The Development of Grain Structure
During Axisymmetric Extrusion of AA3003 Aluminum Alloys. Met. Mater. Trans. A 2015,46, 5920–5932.
26.
Chen, J.Q.; Poole, W.J.; Parson, N.C. The effect of through thickness texture variation on the anisotropic
mechanical response of an extruded Al-Mn-Fe-Si alloy. Mater. Sci. Eng. A 2018,730, 24–35.
27.
Kaneko, S.; Murakami, K.; Sakai, T. Effect of the extrusion conditions on microstructure evolution of the
extruded Al-Mg-Si-Cu alloy rods. Mater. Sci. Eng. A 2009,500, 8–15.
28.
Roy, S.; Allard, L.F.; Rodriguez, A.; Porter, W.D.; Shyam, A. Comparative Evaluation of Cast Aluminum
Alloys for Automotive Cylinder Heads: Part II-Mechanical and Thermal Properties. Met. Mater. Trans. A
2017,48A, 2543–2562.
29.
Nie, J.F.; Muddle, B.C. Microstructural design of high-strength aluminum alloys. J. Phase Equilib.
1998
,
19, 543–551.
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